Laser shock peening apparatuses and methods

ABSTRACT

Methods and apparatuses for processing materials to enhancing the material&#39;s surface strength, improving the material&#39;s cyclic and thermal stability of microstructures, and extend the material&#39;s fatigue performance. Embodiments include laser shock peening at material temperatures that are moderately elevated (from the material&#39;s perspective) above room temperature. Alternate embodiments include laser shock peening at very cold (cryogenic) material temperatures. Still further embodiments include laser shock peening while covering the surface of the material being processed with an active agent that interacts with the laser energy and enhances the pressure exerted on the surface.

This application claims the benefit of U.S. Provisional Application Nos.61/587,815, filed Jan. 18, 2012, and 61/653,050, filed May 30, 2012, theentireties of which are hereby incorporated herein by reference.

FIELD

Embodiments of this disclosure relate generally to changing the surfacestructure and/or properties of a material, and more particularly toshock peening.

Alternate embodiments of the present disclosure relate to materialsurface processing techniques via high power energy pulsed laser shockpeening, and more particularly to methods for enhancing the materialsurface strength, improving the cyclic and thermal stability ofmicrostructures, and extending material's fatigue performance byintroducing dynamic precipitation and dynamic strain aging effectsthrough thermal-mechanical treatments.

BACKGROUND

Conventional material surface processing techniques, such as laser shockpeening (LSP), shot peening (SP) and deep rolling (DR), can be used toimprove fatigue performance of metallic components by generating anear-surface work hardened layer and introducing near-surfacecompressive residual stresses, which retards fatigue-crack initiationand propagation.

Laser shock peening (LSP) is a process that utilizes a high energy pulselaser to induce helpful compressive residual stress and work hardeningsurface layer to a component that results in improved component fatigueperformance and corrosion resistance. A high energy pulse laser can beused to penetrate through the confinement media (water or glass) andshoot onto an ablative coating material (in some embodiments, a metallicthin foil), which vaporizes into a high pressure plasma. The expansionof the plasma generates shock wave propagation into the target componentand plastically deform the component, which generates compressiveresidual stress and work hardening in the component surface.

SUMMARY

Known deformation processing techniques can improve material strength bywork hardening and grain refinement. However, some of these processesinevitably lead to loss of material ductility. Efforts have been made toimprove material strength without severe loss of material ductility.However, some of these methods are not suitable for bulk material and/ornot applicable for industrial manufacturing.

As a severe plastic deformation processing, LSP has limitations. Forexample, the defect density accumulated during LSP can reach saturationquickly and this can limit the material strength improvement throughLSP. By plastic deformation at the cryogenic temperature, dynamicrecovery of dislocations can be suppressed and the saturationdislocation density increased. In this way, the material strength can befurther improved.

Still further, known fatigue strengthening mechanisms do not alwaysremain stable in the presence of cyclic and/or thermal loading. It wasdiscovered that work-hardening and residual stress generated byconventional surface processing techniques can be reduced significantlyduring cyclic loading, particularly at high testing temperature due tothe rearrangement and recovery of microstructures. In this way, theeffect of fatigue life improvement by LSP, SP and DR was limited.Therefore, it was realized that there was a need for better fatigueperformance, which could be realized by stabilizing the microstructuresafter processing.

Embodiments of the present disclosure provide an improved laser shockpeening apparatuses and methods.

Embodiments of the present disclosure provide apparatuses and methodsthat perform laser shock peening while maintaining the material beingprocessed at temperatures that are moderately elevated above roomtemperature and below the melting or recrystallization temperature ofthe material being processed.

Embodiments of the present disclosure provide apparatuses and methodsthat perform laser shock peening while maintaining the material beingprocessed at temperatures that are moderately elevated above roomtemperature and below the recrystallization temperature of the materialbeing processed and generate high dense nano-precipitation for bettermechanical performance, such as fatigue life.

Embodiments of the present disclosure provide apparatuses and methodsthat perform laser shock peening at cryogenic temperatures.

One object of embodiments of the present disclosure is to design ahybrid surface processing technique that integrates laser shock peeningand cryogenic plastic deformation for improved material properties.

An object of alternate embodiments of the present disclosure is todesign a hybrid surface processing technique that perform laser shockpeening at cryogenic temperature to generate deformation twinning andstacking fault for improved mechanical properties.

Embodiments of the present disclosure provide apparatuses and methodsthat perform laser shock peening while covering the surface of thematerial being processed with an active agent that enhances the pressureexerted on the surface.

Objectives of embodiments of the present disclosure include LSPapparatuses and processes that improve material properties, including,strength, ductility and fatigue resistance.

Embodiments of the present disclosure improve upon and resolve thestability problem of current surface processing techniques.

This summary is provided to introduce a selection of the concepts thatare described in further detail in the detailed description and drawingscontained herein. This summary is not intended to identify any primaryor essential features of the claimed subject matter. Some or all of thedescribed features may be present in the corresponding independent ordependent claims, but should not be construed to be a limitation unlessexpressly recited in a particular claim. Each embodiment describedherein is not necessarily intended to address every object describedherein, and each embodiment does not necessarily include each featuredescribed. Other forms, embodiments, objects, advantages, benefits,features, and aspects of the present disclosure will become apparent toone of skill in the art from the detailed description and drawingscontained herein. Moreover, the various apparatuses and methodsdescribed in this summary section, as well as elsewhere in thisapplication, can be expressed as a large number of differentcombinations and subcombinations. All such useful, novel, and inventivecombinations and subcombinations are contemplated herein, it beingrecognized that the explicit expression of each of these combinations isunnecessary.

BRIEF DESCRIPTION OF THE DRAWINGS

Some of the figures shown herein may include dimensions or may have beencreated from scaled drawings. However, such dimensions, or the relativescaling within a figure, are by way of example, and not to be construedas limiting.

FIG. 1.1 —The comparison of surface hardness and plastic deformationdepth after LSP in H₂O₂ and water according to one embodiment of thepresent disclosure.

FIG. 1.2 —Schematic illustration of fast chemical etching-assistedpulsed laser ablation according to one embodiment of the presentdisclosure.

FIG. 1.3 —Mechanism of fast chemical etching-assisted pulsed laserablation according to one embodiment of the present disclosure.

FIG. 1.4 —Ablation depth and volume of PLA with H₂O₂ and H₂Oconfinements according to one embodiment of the present disclosure.

FIG. 2.1 —Schematic representation of a CLSP apparatus according to oneembodiment of the present disclosure. A laser machine is used to deliverthe pulsed laser. Optical mirrors (three in the illustrated embodiment)are optionally used to direct the laser beam. A focus lens can be usedto focus the laser beam to desired beam size. A fixture is typicallyused to fix the target component. The fixture may be placed in a tank(e.g., a stainless steel tank) which can contain a low temperatureliquid.

FIG. 2.2 a —Side view of the CLSP fixture of FIG. 2.1 .

FIG. 2.2 b —Top view of the CLSP fixture of FIG. 2.1 .

FIG. 2.3 —Schematic representation of a gradient nanostructure generatedby CLSP according to one embodiment of the present disclosure.

FIG. 2.4 —TEM image of the deformation twins in stainless steel 304generated by CLSP according to one embodiment of the present disclosure.

FIG. 3.1 —Schematic of a laser beam delivery and focusing system inaccordance with one embodiment of the present invention.

FIG. 3.2 —Schematic of experimental set up and methodology of WLSPaccording to one embodiment of the present disclosure.

FIG. 3.3 —Heating methods for WLSP experimental set up: (A) laserheating; (B) hot gas heating; (C) electric heating; (D) inductionheating according to alternate embodiments of the present disclosure.

FIG. 3.4 —TEM images of microstructures (including highly densenano-precipitate, and dislocation/precipitate entanglement) in aluminumalloy 6061 after WLSP according to one embodiment of the presentdisclosure.

FIG. 3.5 —Extension of fatigue performance of AISI 4140 steel after WLSP(comparing with LSP and non-treated samples) according to one embodimentof the present disclosure.

FIG. 3.6 —The stability of surface hardness during cyclic loading at1400 MPa maximal bending stress of non treated, WLSP and 2 h post-shocktempered samples according to one embodiment of the present disclosure.

FIG. 3.7 —Residual stress relaxation during cyclic loading at 1400 MPamaximal bending stress of WLSP and 2 h post-shock tempered samplesaccording to one embodiment of the present disclosure.

FIG. 3.8 —S—N curves of WLSP and post-shock tempered samples accordingto one embodiment of the present disclosure.

FIG. 4.1 —Surface hardness after WLSP at various laser power intensityand the corresponding peak plasma pressure according to one embodimentof the present disclosure.

FIG. 4.2 —Surface hardness affected by pre- and post-shock tempering atvarious tempering time (laser intensity 3 GW/cm², tempering temperature450° C.) according to one embodiment of the present disclosure.

FIG. 4.3 —Surface hardness affected by the post-shock temperingtemperature at various tempering time (laser intensity 3 GW/cm²)according to one embodiment of the present disclosure.

FIG. 4.4 —TEM images show the initial microstructures of quenched 4140steel without peening according to one embodiment of the presentdisclosure.

FIG. 4.5 —Microstructures in 4140 Steel after WLSP: weak beam dark fieldimages (a and b) and related diffraction pattern of two beam condition(c) showing the entanglement of dislocations and the spherical-shapednano-precipitates (pointed out by arrows) according to at least oneembodiment of the present disclosure.

FIG. 4.6 —Microstructures of WLSP+2 h post-shock tempered sample: (a)bright field TEM image showing dislocation structures; (b) diffractionpattern with weak spots (circled and indicated by the ring) associatedwith precipitate structures; (c and d) dark field TEM images atdifferent magnification showing precipitate structures according to atleast one embodiment of the present disclosure.

FIG. 4.7 —Microstructures in WLSP+6 h post-tempered sample: (a and b)bright field TEM images with different magnification showing dislocationwalls (indicated by arrows) and precipitates with a relatively largesize (circled by rings) are formed and tangled with each other; darkfield images showing precipitates at top surface (c) and a deeper region(d) by selecting the weak spot (circled by the ring) from the inserteddiffraction pattern according to at least one embodiment of the presentdisclosure.

FIG. 4.8 —The evolution of precipitate radius, number density and volumefraction with the post-shock tempering time according to one embodimentof the present disclosure.

FIG. 4.9 —Residual stress relaxation during cyclic loading at 1400 MPamaximal bending stress of WLSP and 2 h post-shock tempered samples. (Theerror bar for the residual stress measurements is 20 MPa) according toone embodiment of the present disclosure.

FIG. 4.10 —The stability of surface hardness during cyclic loading at1400 MPa maximal bending stress of non treated, WLSP and 2 h post-shocktempered samples. The error bar for the hardness test is 10VHN accordingto one embodiment of the present disclosure.

FIG. 4.11 —S—N curves of WLSP and post-shock tempered samples accordingto one embodiment of the present disclosure.

DETAILED DESCRIPTION OF THE ILLUSTRATED EMBODIMENTS

For the purposes of promoting an understanding of the principles of thedisclosure, reference will now be made to one or more embodimentsillustrated in the drawings and specific language will be used todescribe the same. It will nevertheless be understood that no limitationof the scope of the disclosure is thereby intended; any alterations andfurther modifications of the described or illustrated embodiments, andany further applications of the principles of the disclosure asillustrated herein are contemplated as would normally occur to oneskilled in the art to which the disclosure relates. At least oneembodiment of the disclosure is shown in great detail, although it willbe apparent to those skilled in the relevant art that some features orsome combinations of features may not be shown for the sake of clarity.

Any reference to “invention” within this document is a reference to anembodiment of a family of inventions, with no single embodimentincluding features that are necessarily included in all embodiments,unless otherwise stated. Furthermore, although there may be referencesto “advantages” provided by some embodiments, other embodiments may notinclude those same advantages, or may include different advantages. Anyadvantages described herein are not to be construed as limiting to anyof the claims.

Specific quantities (spatial dimensions, temperatures, pressures, times,force, resistance, current, voltage, concentrations, wavelengths,frequencies, heat transfer coefficients, dimensionless parameters, etc.)may be used explicitly or implicitly herein, such specific quantitiesare presented as examples only and are approximate values unlessotherwise indicated. Discussions pertaining to specific compositions ofmatter, if present, are presented as examples only and do not limit theapplicability of other compositions of matter, especially othercompositions of matter with similar properties, unless otherwiseindicated.

Enhanced Laser Shock with Active Liquid Confinement (e.g., HydrogenPeroxide)

Embodiments of the present disclosure include apparatuses and processesto generate an enhanced laser shock with a higher pressure by utilizingactive liquid confinement—hydrogen peroxide (H₂O₂). A mechanism of fastchemical etching-assisted laser ablation is also presented. As a result,the efficiency of underwater laser shock peening of aluminum alloy 6061is improved by 150%, and the ablation rate of pulse laser ablation ofzinc is enhanced by 300%. These methods break a major limitation ofunderwater pulsed laser processing caused by the generation of breakdownplasma.

The laser shock induced by pulse laser ablation under a confinement haspotential for industrial applications, such as laser shock peening(LSP), laser dynamic forming (LDF) and laser-assisted micromachining.The capability, efficiency and application range of these laser-basedtechniques are generally determined by the intensity of laser shock. Theintensity of the laser shock is generally determined by not only thelaser power but also the confining media. The selection of confiningmedia can bring a major limitation to determining the intensity of thelaser shock. For example, solid confinements like glass are inconvenientfor processing 3 dimensional surfaces, and they are also too brittle tostand a high powered laser pulse. On the other hand, even if the liquidconfinements like water are relatively more flexible and practical, theycan still suffer from drawbacks, such as: when the laser intensity isabove a threshold, the breakdown of liquid confinements screens theincident laser power and results in a saturation of laser shock and peakpressure.

Embodiments of the present disclosure include processes and methods toenhance the intensity of laser shock by applying active liquidconfinements like H₂O₂ and other similar compounds. The material to behardened is covered in an active agent (e.g., H₂O₂). The active agentinteracts with the laser energy and enhances the shock peening process.The laser shock is dependent on the high density of laser-inducedplasma, which is determined by the ablation rate of target materialsthrough laser/material interactions. Thus, the intensity of the lasershock can be enhanced by using active liquid confinements because of themore effective ablation caused by the fast chemical etching reactiontaking place simultaneously during the laser ablation process.Experimental results reflect that the surface hardness and plasticdeformation depth of aluminum alloy 6061 (AA6061) after LSP areeffectively increased by using H₂O₂ as the confinement instead of H₂O.LSP efficiency is improved by 150% with the application of H₂O₂. Themechanism described suggests that the enhanced intensity of laser shockproduced by H₂O₂ may be due to the higher ablation rate caused by themutual promotion between the laser ablation and chemical etching. Thismechanism has been further verified by experiment results of pulse laserablation (PLA) of zinc, which reveals that H₂O₂ could dramaticallyincrease the ablation rate by almost 300%. The enhanced laser shock canprovide a great potential in many laser applications.

To verify the enhancement of laser shock by active confinements, LSP ofAA6061 was performed with H₂O₂ as the confinement and aluminum foil asthe ablator, and compared with underwater LSP. H₂O₂ was employed as theactive liquid confinement in this study. H₂O₂ has a strong oxidbillity,low price, and can be applied in industrial applications. A Surelite IIIQ-switched Nd-YAG laser (Continuum Inc.), operating at a wavelength of1064 nanometer with a pulse width (full width at half maximum (FWHM)) of5 nanoseconds was used to deliver laser pulses. The micro-hardness ofsamples after peening was measured by a Leco micro-hardness test machinewith 25 g load and 10s holding time. The surface profile afterprocessing was characterized by a Wyko NT3300 HD surface profiler fromVeeco Inc.

$\begin{matrix}{\varepsilon_{p} = {\frac{2{HEL}}{{3\lambda} + {2\mu}}( {\frac{P}{HEL} - 1} )}} &  {1a} )\end{matrix}$ $\begin{matrix}{{HEL} = {\frac{\lambda + {2\mu}}{2\mu}\sigma_{Y}^{dyn}}} &  {1b} )\end{matrix}$ $\begin{matrix}{{\lambda = \frac{Ev}{( {1 + v} )( {1 - {2v}} )}},{\mu = \frac{E}{2( {1 + v} )}}} &  {1c} )\end{matrix}$

The surface hardness and surface plastic deformation depth after LSPconfined by H₂O₂ and water were measured and compared to evaluate thepressure effect caused by the enhanced laser shock, since the lasershock pressure is relatively difficult to be directly measured. In LSPprocess, the surface plastic strain, ε_(p), depends on the peak pressureof laser shock could be expressed by Eq. 1, where P is the peakpressure, HEL is the Hugoniot elastic limit, σ_(Y) ^(dyn) is the dynamicyield strength, λ and μ are the Lame's constants in terms of Young'smodulus E and Poisson's ratio v. The higher surface strain induced bythe enhanced laser shock pressure brings a stronger strain hardeningeffect, and results in a greater value of surface hardness. On the otherhand, the plastic strain induced by laser shock pressure is a decreasingfunction of the surface depth. The plastic deformation occurs to thedepth where the peak stress no longer exceeds HEL and ε_(p)=0.Therefore, the enhanced laser shock pressure could be characterized bythe surface hardness and plastic deformation depth after LSP.

Experiment results of under liquid LSP are presented in FIG. 1.1 , andeach data point is the average of five measurements. It was determinedthat LSP under H₂O₂ resulted in a greater surface hardness thanunderwater LSP (FIG. 1.1A). For instance, with the same laser intensity8 GW/cm², the hardness was increased from 113 VHN (Vickers hardnessnumber) to 123 VHN by applying H₂O₂. Considering that the hardness ofAA6061 before LSP is around 94VHN, H₂O₂ increases the efficiency ofunderwater LSP by 153%. In addition, it is also found that even if thelaser intensity of underwater LSP is enhanced from 4 GW/cm² to 12GW/cm², the hardness is increased by 9% from 108VHN to 118VHN. This islikely because while the laser intensity is over 10 GW/cm², thebreakdown plasma is generated, and results in the saturation of peakpressure. On the other hand, replacing H₂O by H₂O₂ can break thislimitation and enhance the hardness up to around 132VHN, which is closeto the saturated hardness value of AA6061 caused by the strainhardening.

Meanwhile, H₂O₂ brings a relatively larger surface deformation depth aswell (FIG. 1.1B). For example, with the same laser intensity 10 GW/cm²,H₂O₂ increased the deformation depth by 153.4% from 0.73 μm to 1.12 μm.Furthermore, the same saturation phenomenon of laser shock was observedfor underwater LSP, and H₂O₂ is believed to effectively break thislimitation by forming a deeper surface dent. Therefore, both surfacehardness and deformation depth after LSP suggest that the enhanced lasershock with a greater pressure is generated by applying H₂O₂ as theconfining media.

The intensity of laser shock is likely introduced and governed by theexpansion of the laser-induced plasma, while the plasma parametersincluding density, pressure and temperature are ruled over by theablation rate. Therefore, it is believed that the greater ablation ratepromoted by active confinements may play a role in this process toenhance the laser shock with a higher pressure.

To help understand this process, the mechanism of fast chemicaletching-assisted PLA was applied, which is schematically illustrated inFIG. 1.2 . The flow chart of this mechanism is presented in FIG. 1.3 ,and the ionization and chemical etching reactions taking place duringthe process are summarized as Eq. 2-6 (see FIG. 1.3 ), where hv standsfor the photon energy, and ΔE_(T) is the thermal energy released byetching reactions. When the surface of metal target at the focal spot isvaporized and ionized by the front part of incident laser pulses, alaser-induced plasma plume is created at the target/confinementinterface (FIG. 1.2A). This ionization is governed by theinverse-bremsstrahlung mechanism (Eq. 2, see FIG. 1.3 ) related to thefast growth of free electrons e⁻. Once generated, the plasma isimmediately forced into a thermodynamic state due to the confiningeffect of the liquid confinement. Meanwhile, a dramatic increase of theplasma pressure and temperature is induced by absorbing the later partof laser energy, followed by the formation of a shockwave with asupersonic velocity at the interface. Many reports have alreadydemonstrated that the pressure of the laser-induced plasma under liquidconfinements is at the level of GPa, and the temperature could attainseveral thousand K. This environment provides for the decompositionreaction of H₂O₂, which releases the atomic oxygen O⁻ as a strongoxidizer for the chemical etching process (Eq. 3). (Note that thebyproduct of this process using hydrogen peroxide as the active agent iswater, which is easily disposed of, recycled, or reprocessed back tohydrogen peroxide). Noting here, the high temperature could acceleratethe chemical decomposition rate. The decomposition rate of H₂O₂ within ahigh temperature environment (higher than 10³K) can reach the order ofmagnitude of 10⁷ cm³ mol⁻¹ s⁻¹. In addition, considering the powerfulcondensation effect of liquid confinements, it is reasonable to indicatethat the decomposition reaction occurs at the surface region of theplasma plume and the released atomic oxygen is confined and exists atthe interface as illustrated in FIG. 1.2B. Furthermore, some releasedatomic oxygen could be further ionized by absorbing the later laserenergy to form ionized oxygen O⁺ with an even stronger oxidbillity dueto the inverse-bremsstrahlung process (Eq. 4, see FIG. 1.3 ). Thus, thefast chemical decomposition rate and small confining volume could causethe formation of a plasma thin layer of atomic and ionized oxygen with astrong oxidbillity and high concentration, resulting in a fast chemicaletching process to remove more target materials (Eq. 5, 6, see FIG. 1.3).

The mutual promotion of the greater ablation rate and enhanced plasmaresults in a further increase of ablation rate, and thus the laser shockpressure. The higher plasma pressure results in a higher ablation rate.According to the discussion above, the enhanced plasma with a higherinternal energy and pressure is expected due to the following reasons.First of all, the chemical etching process increases the plasma densityand thickness by removing more target materials, resulting in anincrease of plasma pressure. The total pressure of plasma P is the sumof the electron partial pressure P_(e) and the particle partial pressureP_(p) as illustrated by Eq. 7, where k_(B) is the Boltzmann constant,n_(e) and n_(p) are the amount of electron and particle relatively.Therefore, increasing electron and particle density n_(e) and n_(p) bychemical etching reactions could strengthen the plasma pressure.Furthermore, based on Fabbro's model, the time evolution of the plasmapressure P(t) and thickness L(t) is presented in Eq. 8, where Z₁ and Z₂are the shock impedances of target and confining media. It isdemonstrated the plasma pressure is proportional to the plasmathickness. The greater plasma thickness is expected, since the greaterablation rate could cause a deeper ablation depth. Secondly, the plasmaparameters could be further affected by the thermal energy released fromchemical etching process (Eq. 5, 6). This thermal energy will be appliedto increase the plasma internal thermal energy and open theliquid/target interface. The energy balance during processing could bedescribed by Eq. 9, where I, E_(T), and α are the laser intensity,thermal energy absorbed by plasma, and a constant fraction respectively.Eq. 9 shows that increase of released thermal energy ΔE_(T) could causea higher plasma pressure. Thus, the ablation rate and plasma pressurereciprocally promote each other, and result in an enhanced laser shock.

$\begin{matrix}{P = {{P_{e} + P_{p}} = {{k_{B}T_{e}n_{e}} + {k_{B}T_{p}n_{p}}}}} &  7 )\end{matrix}$ $\begin{matrix}{{\frac{\mathbb{d}{L(t)}}{\mathbb{d}t} = \frac{2{P(t)}}{Z}},{{2/Z} = {{1/Z_{1}} + {1/Z_{2}}}}} &  8 )\end{matrix}$ $\begin{matrix}{{\frac{\mathbb{d}( {\Delta E_{T}} )}{\mathbb{d}t} + {I(t)}} = {{{{P(t)}\frac{\mathbb{d}L}{\mathbb{d}t}} + \frac{\mathbb{d}\lbrack {{E_{T}(t)}L} \rbrack}{\alpha{\mathbb{d}t}}} = {{{P(t)}\frac{\mathbb{d}L}{\mathbb{d}t}} + {\frac{3}{2a}\frac{d\lbrack {{P(t)}{L(t)}} \rbrack}{dt}}}}} &  9 )\end{matrix}$

To verify the mechanism that enhanced laser shock is caused by thegreater ablation rate induced by active confinements, PLA of Zinc wascarried out with H₂O₂ as a confinement medium and compared with PLA ofZinc being carried out with H₂O as a confinement medium. The increase ofablation volume by the active confinement H₂O₂ is demonstrated in FIG.1.4 . The cross-sectional profiles of the ablated area after PLA areshown in FIG. 1.4A, and an image of ablated area taken by the profileris inserted at the lower left corner. To make the differences measurableand reduce the error introduced by the laser pulse variance, eachprofile was obtained with about 5 pulses PLA with a laser fluence ofabout 30 J/cm². It is observed that the maximum depth at the center ofthe ablated area was dramatically increased by approximately 200% fromabout 3 μm to about 9 μm while applying H₂O₂ as the confinement insteadof H₂O. This is evidence of increased plasma thickness in the mechanism.In addition, the ablation diameter (R and R′) was enhanced by about 16%from about 0.32 mm to about 0.37 mm as well. This indicates theformation of the plasma thin layer of atomic and ionized oxygen asdiscussed above. The ablation volume V_(A) was also compared to theablation rate. FIG. 1.4B illustrates the increase of ablation volume byH₂O₂, and each data point is the average of 5 measurements. Forinstance, with 5 pulses ablation at a laser fluence of 60 J/cm², theablation volume was dramatically enhanced by about 301% from about3.39E-6 mm³ to about 1.36 E-5 mm³; with 10 pulses, the ablation volumewas enhanced by about 321% from about 1.05E-5 mm³ to about 4.42E-5 mm³.Furthermore, it was also found that when using water as the confinement,even if the laser fluence is increased from 30 J/cm² to 60 J/cm², theablation volume was slightly affected, such as enhanced by about 25.9%from about 8.34E-6 mm³ to about 1.05E-5 mm³ with 10 pulses ablation.This may be explained by the generation of breakdown plasma, whichscreens the laser pulse from reaching the target. However, with the samelaser fluence about 30 J/cm² of 10 pulses ablation, and replacing H₂O byH₂O₂ as the confinement, can dramatically enhance the ablation volume byabout 143.4% from about 8.34E-6 mm³ to about 2.03E-5 mm³. Applyingactive confinement has a much higher efficiency on enhancing ablationrate than increasing laser power. This benefit can be broadly applied inpulsed laser processing, such as saving laser power and enhancingworking efficiency in laser micromachining, and increasing theproductivity in laser-assisted nanocrystal synthesis.

This work presents new apparatuses and processes to generate theenhanced laser shock with a higher pressure by applying H₂O₂ and othersimilar compounds as an active liquid confinement. The mechanism of fastchemical etching-assisted laser ablation is now possible using enhancedlaser shock produced by H₂O₂ as a confinement medium. The greaterablation rate induced by the mutual promotion between ablation andetching processes is believed as the key factors for the enhanced lasershock. As a result, this method breaks a current limitation inunderwater pulsed laser processing caused by the generation of breakdownplasma, and effectively improves the efficiencies of underwater LSP andPLA by about 150% and about 300%, respectively. Thus, the enhanced lasershock by active confinement can extend the applications of pulsed laserprocessing, such as LSP, laser-assisted micromachining, and under liquidPLA.

Cryogenic Laser Shock Peening

Hybrid surface processing methods and apparatuses integrating lasershock peening (LSP) and cryogenic plastic deformation is describedaccording to various embodiments of the present disclosure. Bysubmerging the target component into a low temperature liquid duringLSP, helpful microstructure change can be induced in the targetcomponent. By manipulating the processing parameters, gradientnanostructure can be generated in material for best combination ofmaterial strength and ductility.

Cryogenic laser shock peening can complement traditional laser shockpeening process for improved material properties including strength,ductility and fatigue performance. As an example, cryogenic laser shockpeening can induce helpful microstructure change, including higherdefect density, gradient microstructure, etc., in metallic componentsfor enhanced material properties.

Cryogenic laser shock peening (CLSP) systems and methods according toembodiments of the present disclosure include methods and fixtures toimpose low temperature to the target component. During the CLSP process,the target sample was kept at the cryogenic temperature by differentmethods. The temperature was controlled with a cryogenic thermalplatform, such as a cold plate (temperature variable) or a cold liquidtank (temperature fixed). The sample temperature can be manipulated bysubmerging the target samples into the tank. The tank temperature couldbe changed by utilizing different cold liquids such as liquid nitrogen,liquid helium, liquid CO2, water-ice mixture, etc. Adhesive paints oraluminum foils can be used as the ablative coating materials protectingthe samples from surface melting by the pulsed laser. For samples withflat surface, glass can be used as confinement media. Samples withcurved surface will be submerged at least partly into the liquid bath.In this setup, the cold liquid serves as the cooling media as well asthe confinement media. First, the cold liquid cools the samples to theboiling temperature of the cold liquid. Secondly, the cold liquidservers as the confinement media, similar to under water laser shockpeening. Cryogenic systems have been used in industry for purposes suchas cryogenic machining, cryogenic rolling, and cryogenic equal channelangular extrusion. In one embodiment to utilize CLSP in the industry,the tank can be replaced by a nozzle, which will spray cold liquid tocover the target sample both to cool the sample and to serve asconfinement media.

The low temperature during CLSP effectively suppress dynamic recoveryduring the CLSP process and thus significantly improves materialstrength. The other effect of CLSP is the generation of highly densenanomaterials (for example, nano-twins) at material subsurface asillustrated in FIG. 2.4 . It has been reported that nanomaterials, suchas nano-twins, can effectively block dislocation movement like grainboundaries and improve material strength.

In addition, CLSP can generate gradient nanostructure for a usefulcombination of strength and ductility. By controlling the parameters,such as the laser intensity, laser beam size, peening overlap ratio, thedistribution and range of the gradient nanostructure can be controlledfor different materials for a useful combination of materialperformance.

The optimal temperature at which the target material is maintained forCLSP varies with the properties of the target materials. The optimaltemperature will generally maximize the mechanical twining, stackingfault in metals and/or martensite in steel. Methods to find out theoptimal temperature range of CLSP for metals, including copper alloy,steel, and NiTi alloy were discovered. For many materials, the optimaltemperature is from approximately 50 deg. K to approximately 250 deg. K.In general, the flow stress should be higher than the critical twinningstress in order to generate deformation twinning. CLSP can generate highvolume fraction of martensite, high density stacking faults, and highdensity deformation twins resulting in higher material hardness andbetter fatigue performance than other forms of LSP.

FIG. 2.1 depicts a cryogenic laser shock peening method and apparatusaccording to one example embodiment. A high energy pulsed laser wasdelivered by a laser machine. Three optical mirrors were used to directthe laser beam, although alternate embodiments include a differentnumber of mirrors or no mirrors. A focus lens was optionally used tofocus the laser beam. The confinement media, the ablative coatingmaterial and the target component were clamped together by a fixture. Aside-view of the fixture was shown in FIG. 2.2 a . The target componentwas optionally on top of the fixture. On top of the target component wasthe ablative coating material (a metal foil with thickness around 30 to200 μm). On top of the coating material was the confinement media (glassor other transparent solid material). Fasteners (e.g. four screws) wereoptionally used to clamp the confinement media, the ablative coatingmaterial and the target component to the fixture. FIG. 2.2 b shows thetop-view of the fixture.

During peening, a high energy pulsed laser penetrated through theconfinement media and shot onto the ablative coating material. A portionof the ablative material was vaporized and generated high temperatureplasma. The expansion of the plasma generated shock wave propagationinto both the confinement media and the target component. This generatedplastic deformation and helpful compressive residual stress in componentsurface. During the process, the component was cooled by, for example,being submerged in low temperature liquid. This liquid was optionallystored in a stainless tank. Depending on the low temperature, differentliquid, such as liquid helium, liquid nitrogen (LN₂), liquid carbondioxide, could be used. The low temperature of the component effectivelysuppressed dynamic recovery during plastic deformation and thus lead tohigh defect density in the material. In addition, the low temperaturecould effectively increase the density of deformation twins generated invarious metallic materials (copper and its alloys, Iron, carbon steel,stainless, titanium and its alloys, etc.). Like grain boundaries, thesedeformation twins effectively blocked dislocation movement and thushardens the materials. In addition, the deformation twins had goodthermal stability and lead to good thermal stability of themicrostructure after CLSP. Furthermore, CLSP generated higher volumefraction of martensite transformation than at room temperature LSP andthus further strengthen the material.

One feature of the CLSP process was the generation of gradientnanostructure. As shown in FIG. 2.3 , after peening both sizes of atarget component, gradient microstructures were generated. In the topsurface, where some severe plastic deformation occurs,nano-crystallization occurred. The nano-crystalline surface had highstrength and was effective to inhibit crack formation. At materialsubsurface, some plastic deformation took place and the grains wererefined to a certain extent. In this layer, the microstructure wascharacterized by refined grain with dislocations and deformation twinsembedded in the grains. The interaction between the dislocation anddeformation twins effectively strengthened the material and providedgood thermal stability. Deeper into the material, the material was lessaffected by the peening process and the microstructure was characterizedby coarse grains with relatively low density dislocations. Due to itslow dislocation density and large grain size, the materials in thisregion had enough room for strain hardening (work hardening) and thusrendered good ductility of the component during tensile loading.

It was believed that CLSP could effectively induced gradientnanostructure in metallic components for improved material strength andductility. It was noted that various changes or modifications were madeto the process parameters to suit different requirement in real cases.

Embodiments include one or more of the following:

-   -   In the cryogenic laser shock peening process (CLSP), the target        component is kept at the desired low temperature through a low        temperature liquid bath during the peening process.    -   The low temperature can effectively suppress dynamic recovery        and thus lead to high defect density, which further hardens the        material.    -   The low temperature can induce the generation of highly dense        deformation twins in various metallic materials (Cu and its        alloy, iron, carbon steel, stainless steel, etc.). These        deformation twins harden the material like the grain boundaries        and have good thermal stability.    -   CLSP can induce higher volume fraction of martensite        transformation than room temperature LSP and thus further        strengthen the material.    -   CLSP can generate gradient nanostructure in metallic component,        which leads to high material strength and ductility.        Warm Laser Shock Peening and Optional Heat Treatment

Embodiments of the present disclosure utilize warm laser shock peeningand optional heat treatment, which can stabilize residual stresses,increased strength, and enhanced fatigue performance.

Further embodiments of the present disclosure include warm laser shockpeening (WLSP), which can be a rapid and low cost thermal-mechanicalmaterial processing technique for enhancing the fatigue performance ofmetallic components and improving the cyclic and thermal stability ofmicrostructures via high energy pulsed laser shock peening at elevatedtemperatures. During an example process, the laser energy was absorbedby the ablative coating materials placed on the surface of targetmaterial to generate laser-induced plasma. The hydrodynamic expansion oflaser-induced plasma was confined by the presence of a transparentconfinement, leading to the generation of high pressure shock waves thatpropagated into the target and resulted in a near-surface high strainrate plastic deformation. Meanwhile, a heating source was applied toelevate the temperature of target to a specified temperature or range oftemperatures. The results from mechanical processing and thermal energylead to the generation of highly dense nano-scale precipitates anduniformly distributed dislocation structures through dynamicprecipitation (DP) and dynamic strain aging (DSA). As compared toconventional surface processing techniques (such as traditional LSP),WLSP effectively improved the fatigue performance of metalliccomponents. The interaction between nano-precipitates and dislocations,also known as dislocation pinning effect, contributed to the fatigueimprovement through the enhanced cyclic and thermal stability ofmicrostructures. Warm laser shock peening (WLSP) and the followingpost-shock tempering treatment have been found to achieve optimizedsurface strength and fatigue performances of metallic materials. Thisimproved material stability and reliability were attributed to theenhanced dislocation pinning effect corresponding with the numberdensity, size and space distribution of nano-precipitates, which couldbe tailored by manipulating processing conditions of WLSP and post-shocktempering. The precipitation kinetics as well as dislocation pinningstrength affected by precipitate parameters were analyzed. The heatingmethods, such as laser heating, electromagnetic heating, hot gasheating, or electric heating were feasible to heat the metal pieceduring the process. The temperature during laser shock peening washelpful to the final microstructure and mechanical properties. Optimalprocessing temperatures were discovered during WLSP. The post annealingtemperature and time were also useful to control the microstructure foroptimal mechanical properties (see appendix).

The optimal temperature will generally vary based on differences in thetarget material. The optimal temperature will maximize the density ofnanoprecipitation by mechanism of dynamic strain aging and dynamicprecipitation. Normally the optimal temperature is below therecrystallization temperature and higher than the room temperature. Iftemperature is too high, dislocation density will reduce and less amountof nucleation size for nanoprecipitation, ending up with coarseprecipitates. If temperature is too low, the nucleation rate of theprecipitation will be reduced, ending up with very small amount ofnanoprecipitates.

Referring to the example embodiment depicted in FIG. 3.1 , a Q-SwitchedNd-YAG laser was used to deliver the laser pulse. The laser wavelengthwas adjusted as 1064/532/355 nm dependent on the absorption of theablative coating. A suitable laser system of an energy output in therange of several to several hundred J/pulse with a pulse duration ofoptionally less than 100 ns was used. A focus lens was optionallyapplied to focus the laser beam on the surface of target to a certainsize. An optional motorized X-Y stage was placed at the bottom of targetto manipulate the laser scanning speed and laser processing positions.The laser scanning speed was accurately determined by considering thelaser frequency and beam size plus the laser beam overlap ratio.

FIG. 3.2 depicts a schematic of an example experimental set up andmethodology of WLSP according to at least one embodiment. During WLSP, alaser pulse penetrated through the confining media and irradiated theablative coating material. With the absorption of the laser energy, theablative coating material at the focal spot was vaporized and ionized,resulting in the formation of laser-induced plasma plume. The expansionof plasma, which was optionally confined by the presence of confiningmedia, led to the shock wave propagation into the target material andthe optional confining media. When the shock wave pressure exceed theHugoniot elastic limit of target material, plastic deformation occurs,accompanied by the generation of near-surface work hardening layer andcompressive residual stress. Besides the plastic deformation, thethermal energy provided by the heating sources during WLSP playedanother useful role. The combined effects of mechanical deformation andthermal heating lead to the generation of highly dense nano-precipitatesin target materials through dynamic precipitation process. The elasticinteraction between nano-precipitates and nearby mobile dislocations, socalled dislocation pinning effect, stabilized the microstructures duringcyclic and/or thermal loading, and resulted in improved fatigueperformance.

The optional confining media of WLSP can generally be any transparentmaterial, such as glass, water, fused quartz and silicone oil. Note thatfor the industrial applications of WLSP, the silicone oil with a flashpoint over 600° F. is generally more practical as a confining media thanthe glass and water, since the glass could have be shattered by thelaser-induced shock pressure, and the vapor point of water was too low.

The ablative coating materials (metallic foils, organic paints oradhesives) were used for the absorption of laser energy and alsoassisted the surface of target material from melting, ionizing anddamaging by laser pulse.

The laser power intensity was one helpful laser parameter in WLSPprocess, since it was related to the shockwave peak pressure thatgoverned the magnitude and depth of compressive residual stress. Theplastic deformation occurred when the shock wave pressure exceed theHugoniot elastic limit of target material. With the increase of laserpower intensity, residual stresses increase with depth but decrease atthe surface because of surface release waves. Additionally, thedislocation density, which strongly affects the nucleation processthrough dynamic precipitation, reached a saturation point associatedwith a certain magnitude of shockwave peak pressure. The above factorsindicate that there were optimal WLSP conditions in terms of laser powerintensity. The laser power intensity was affected and could be adjustedby several other laser parameters including laser pulse duration, laserbeam size, pulse energy, etc.

Besides laser power intensity, the WLSP processing temperature playedanother useful role, and was accurately manipulated. When the processingtemperature was lower than the optimal one, the dynamic precipitationeffect was less effective, resulting in a weaker dislocation pinningstrength and thus a worse fatigue performance. On the other hand, whenthe processing temperature was too high, the thermal-induced dynamicrecovery occurred, leading to a softening effect on the surfacestrength.

The resulting microstructures and stress/strain distribution after WLSPwere strongly affected by the above two processing parameters: laserpower intensity and processing temperature. Greater laser powerintensity led to a higher strain rate plastic deformation by generatinga stronger shockwave peak pressure, which governed the magnitude anddepth of resulted compressive residual stresses. The relationshipbetween laser power intensity I₀ and peak pressure P was expressed byFabbro's model as

${{P( I_{0} )} = {0.1( \frac{\alpha}{{2\alpha} + 3} )^{1/2}Z^{1/2}I_{0}^{1/2}}},$where α is the efficiency of the interaction. Z is the combined shockimpedance defined as Z=2/(1/Z₁+1/Z₂), where Z₁ and Z₂ were the shockimpedance of the material and the confining media, respectively.Additionally, the strain hardening and dynamic strain aging effects wereenhanced due to stronger laser intensity, leading to the formation ofdislocation structures with a higher density. These dislocationsinteracted and tangled with each other to form the dislocation cores,which were accepted as the available nucleation sites where it waspossible for the nucleation to take place. This was due to thedislocation core energy released for nucleation and the enhancednucleation atoms' migration assisted by the dislocationinter-connections. The laser power intensity increased the nucleationrate of nano-precipitates by generating more available nucleation sitesand providing dislocation core energy. On the other hand, the elevatedprocessing temperature had a favorable influence on the nucleation rateduring WLSP by providing thermal energy for nucleation reactions andalso decreasing the chemical driving force for nucleation. The chemicaldriving force represented the excess free energy provided by theunstable phase relative to the stable phase. A precipitate was generallyconsidered as the relatively stable phase with lower free energy beforereaching the equilibrium temperature, and the excess free energyprovided by the unstable phase (matrix material) was reduced byelevating the temperature. As one major component of nucleationactivation energy, the chemical driving force decreased by the elevatedtemperature noticeably improve the efficiency of WLSP by increasing thenucleation rate.

FIG. 3.3 depicts the heating methods which can be used for WLSP, such aslaser heating, hot gas heating, electric heating, and induction heating.For laser heating systems (FIG. 3.3A), the positions of sample surfacesubjected to laser peening process were simultaneously treated byanother laser beam which provided laser energy to heat up thetemperature of focal points. For hot gas heating systems (FIG. 3.3B),the hot gas with a certain temperature was generated by a hot gasgenerator and inputted into a gas chamber to elevate the targettemperature. For the electric heating equipments such as electricheating plate (FIG. 3.3C), the equipment could be placed at the bottomof the target to elevate the target temperature by converting electricalenergy to thermal energy. For induction heating systems (FIG. 3.3D), apower supply sends an AC current through an inductor (copper coil) toheat up the work piece placed inside the inductor.

The accurate measurement of residual stress was useful in the qualitycontrol of WLSP process. Several techniques could have been used tomeasure residual stress, including center-hole drilling, X-raydiffraction, layer removal, etc. The center-hole drilling methodinvolves drilling a small hole into the sample, and using a residualstress gauge to measure the relieved surface strains. The X-raydiffraction technique uses the lattice spacing as the strain gauge tocalculate the residual stress based on the changes in Bragg angle andinter-planar spacing. The layer removal method involves the removal oflayers on one side of a flat plate, and the resulted curvaturesdependent on the original stress distribution are measured for theresidual stress deduction.

TEM images in FIG. 3.4 show the microstructures of aluminum alloy 6061after WLSP. It was observed that the highly dense spherical-shapednano-precipitates with a diameter around 6-10 nm were generated afterprocessing (pointed out by arrows). The nucleation took place duringWLSP was assisted by the thermal energy supplied by the warm temperatureand the high strain rate plastic deformation provided by the laser shockpressure. As the high strain rate causing plastic deformation occurred,the highly dense dislocations were generated, which interacted andtangled within each other and with nano-precipitates. Consequently,highly dense and uniformly distributed dislocations were formed due tothe particle/dislocation interaction.

FIG. 3.5 depicts the stress-lifespan (S—N) curves for bending fatiguetesting of AISI 4140 steel after various processing conditions. The S—Ncurve moved to the right after LSP and WSLP. Under certain stressmagnitudes (1200 and 1500 Map), the fatigue life after WLSP was 3-5times higher than that after LSP. This notably indicated that WLSPimproved fatigue performance more effectively than LSP. This was mainlyresulted from the enhanced cyclic and thermal stability after WLSP dueto the dislocation pinning effect.

FIG. 3.6 compares the residual stress relaxation during the cyclicloading at 1400 MPa maximal three-point bending stress of non-treatedsamples, samples treated with WLSP, and samples treated with WLSPfollowed by two hours of annealing (2 h post-shock tempered samples).Compared to WLSP samples, the initial residual stress before cyclicloading of 2 h post-shock tempered samples had a lower magnitude. Thisstress relaxation was mainly caused by the thermal-induced diffusionalcreep and dislocation glide during the tempering process. However, amore noticeable stress relaxation could be seen in WLSP samples withincreasing cyclic loading numbers.

In WLSP samples, the cyclic residual stress relaxation was likely due tothe micro-plastic strains accumulated from cycle to cycle, while in 2 htempered samples, the presence of larger size precipitate effectivelyresisted material's plastic behaviors by locking mobile dislocations. Inaddition, the stress relaxation occurred when the superposition of theloading stress, residual stress and dislocation pinning stress exceededa material's yield stress. Therefore, for 2 h tempered sample, with thesame loading stress and a lower initial residual stress, the suppressedstress relaxation was attributed to the enhanced pinning stress. Thisimprovement made a major contribution for a better fatigue performance.

The stability of surface strength was also useful for the effectivefatigue improvements by determining the resistance to crack initiation.The surface hardness changed with the cyclic loading of samplesprocessed by various conditions is shown in FIG. 3.7 . It was ofinterest to notice that in the first 100 cycles, a helpful hardeningeffect was induced by cyclic loading in 2 h tempered samples, while nosuch phenomenon was observed in WLSP samples. Both WLSP and 2 h temperedsamples contained the highly dense dislocations, but the precipitatesize of 2 h tempered sample was 4 times larger than that of WLSP sample.Therefore, the cyclic hardening in 2 h tempered sample was explained bythe enhanced pinning effect from the precipitate coarsening. Theresulting larger particles exert a greater pinning force to resist themovement of dislocations during the cyclic loading. To continue theplastic behavior, the generation of more mobile dislocations wasnecessary, resulting in the enhanced dislocation multiplication. Thiscontributed to a greater surface hardness. After 100 cycle loadings, asoftening effect of surface hardness was observed with increasing cyclenumbers in both WLSP and 2 h tempered samples. However, compared to theWLSP sample without additional hardening, the 2 h tempered sample had alower softening rate, which indicated a higher cyclic stability ofsurface strength.

Three-point bending fatigue performances obtained after variousprocessing conditions were demonstrated by the stress-lifespan (S—N)curves in FIG. 3.8 . It was shown that compared to WLSP, a betterfatigue performance was achieved by 2 h post-shock tempering treatment.

Thermal Engineered Laser Shock Peening

Embodiments of the present disclosure utilize thermal engineered lasershock peening, which can enhance fatigue performance.

Thermal engineered laser shock peening (LSP) according to embodiments ofthe present disclosure is a technique combining warm laser shock peening(WLSP) with the followed post-shock tempering treatment to achieve theoptimized surface strength and fatigue performances of metallicmaterials. This technique integrates aspects of LSP, dynamic strainaging (DSA), dynamic precipitation (DP) and post shock tempering toobtain optimized microstructures for extending the fatigue life, such asnano-precipitates and highly dense dislocations. In this work, AISI 4140steel was used to evaluate thermal engineered LSP process. The resultingmicrostructures as well as mechanical properties were studied undervarious processing conditions. The mechanism of fatigue performanceimprovements was investigated. It was found that the extended fatiguelife was generally caused by the enhanced cyclic stability ofcompressive residual stress as well as surface strength. This improvedmaterial stability and reliability were attributed to the enhanceddislocation pinning effect corresponding with the number density, sizeand space distribution of nano-precipitates, which could be tailored bymanipulating processing conditions of WLSP and post-shock tempering.Precipitation kinetics as well as dislocation pinning strength affectedby precipitate parameters were analyzed.

Laser shock peening (LSP) is a surface processing technique used toextend the fatigue life of metallic materials. The near-surfacecompressive residual stresses as well as work-hardening states arecharacteristics produced after LSP, resulting in the high resistance tothe crack initiation and propagation. These near-surface alterationsmake contributions for the improvement of fatigue performance. Due tothe high strain rate deformation caused by the laser shock pressure(normally 10⁵˜10⁷/s), the magnitude and depth of compressive residualstresses as well as the strain-hardening ratio produced by LSP aregenerally greater than other plastic deformation processes such as shotpeening and deep rolling. However, the compressive residual stress maynot be stable during the cyclic loading, especially under the hightesting temperature. The residual stress relaxation of laser shockpeened steel AISI 304 was increased with increasing the temperature andcyclic loading number. The similar phenomenon was reported that both theannealing treatment and the cyclic loading lead to helpful relaxation ofsurface residual stress of steel 4140 processed by LSP. This stressrelaxation related to the microstructure rearrangement can make LSP lesseffective for fatigue improvement.

Warm laser shock peening (WLSP) can increase the stability ofmicrostructures. In some embodiments, WLSP is a thermal-mechanicalprocessing technique, which integrates the aspects of LSP, dynamicstrain aging (DSA), and dynamic precipitation (DP) to enhance fatigueperformance. DSA was a strengthening mechanism caused by the interactionbetween the mobile dislocations and diffused solute atoms. Thisinteraction could effectively enhance the dislocation multiplication,which lead to highly dense dislocations with a uniformly distributedarrangement. DP, also known as strain-induced precipitation, took placesimultaneously during the deformation to generate highly dense nanoscaleprecipitates. The nucleation process during DP assisted and acceleratedby the presence of highly dense dislocations from DSA. The numberdensity of nano-precipitates generated by WLSP is in the order ofmagnitude of 10⁴/μm³. These precipitates act as local barriers andinteract elastically with nearby moving dislocations by exerting apinning force to inhibit and resist dislocation movements. Thisprecipitate/dislocation interaction, so-called the dislocation pinningeffect, effectively contributed to the stabilized microstructures, whichin turn improved the stability of surface strength and compressiveresidual stress and thus extends the fatigue life. For example, furtherimprovements for the surface hardness of AISI 1042 steel was produced byWLSP relative to LSP. Compared to LSP, WLSP results in a better fatigueperformance of 4140 steel due to the improved residual stress stability.However, the improved surface hardness of 1042 steel by WLSP relative toLSP was less than 10%; while the fatigue limit of 4140 steel after WLSPis increased by 75 MPa corresponding to 7%. Therefore, the efficiency ofWLSP on mechanical performances was low, and the microstructures afterWLSP was not optimized. It was useful to discover a new approach thatfurther improved WLSP efficiency so as to obtain the optimalmicrostructures for mechanical applications.

From the point of view of dislocation pinning strength, the limitationof WLSP efficiency was generally attributed to the limited pinningstrength due to the precipitate size (normally 5-10 nm in diameter).This was determined by the short DP time (in nanosecond level)associated with the laser pulse duration. Pinning strength is stronglyaffected by several precipitate parameters including the size, numberdensity, inter-particle spacing, and volume fraction. To overcome theprecipitate/dislocation interaction, a greater applied stress was neededfor the metal matrix containing particles with a higher volume fraction,larger size and/or greater number density. The dislocation pinningeffect induced by WLSP had been studied by multiscale discretedislocation dynamic (MDDD) simulation. It was found that with theconsideration of precipitate volume fraction, the optimal pinningstrength was determined by balancing the particle size and numberdensity effects. MDDD simulation results indicate the microstructuresafter WLSP are not necessarily optimized for mechanical performances,and have a potential to be further improved.

In order to obtain the optimal microstructures with the optimizedpinning strength for mechanical applications, thermal engineered LSP wasstudied. This technique extended the precipitation kinetics from thenucleation stage to the coarsening stage by combining WLSP with thefollowed post-shock tempering treatment. Tempering was a heating processused to treat the quench-hardened steel for a combination of strength,ductility and structural stability. The coarsening of nano-precipitateduring the post-shock tempering differs from the conventional staticaging due to the presence of interconnected dislocation network fromDSA, which could effectively assist and accelerate the solute diffusion.Cyclic loading induced the formation of precipitates in Al—Mg—Simaterials due to the presence of highly dense dislocations. A numericalmodel was created to study the precipitation kinetic in Nb microalloyedsteel by considering the heterogeneous nucleation on dislocations andthe enhanced growth due to the pipe diffusion through the dislocationnetwork. During thermal engineered LSP, the enhanced pinning effect wasachieved by a short time tempering treatment after WLSP that providedthe thermal energy and the aging time for the growth of precipitates.However, this dislocation-assisted coarsening of particles was mostlyaccompanied with a decrease of particle number density. Theover-tempering effect weakened the pinning strength by decreasing theparticle number density, and lead to the dislocation thermal glide andresidual stress relaxation. As a result, the post-shock temperingtreatment was investigated by carefully manipulating tempering time andtemperature to control the balance between particle size and numberdensity for an optimal pinning effect. In addition, thermal engineeredLSP was studied for the optimal processing condition, since it wasthought that this technique had a greater potential than other surfaceprocessing techniques (shot peening, deep rolling, LSP etc) to reach theoptimal mechanical performances of metal materials due to: (1) theeffective strain hardening effect and the high magnitude and in-depthcompressive residual stress produced by the ultra high strain rateplastic deformation during LSP; (2) the highly dense dislocations with auniform arrangement induced by DSA; (3) the high nucleation rate ofnano-precipitates through DP (10⁴/μm³ [8]); and (4) the optimizedpinning strength obtained by the post-shock tempering treatment.

To study these concepts, thermal engineered LSP was carried out on AISI4140. The microstructures after processing with various processingconditions were characterized by TEM. The resulted mechanical propertiesincluding surface hardness, cyclic stability of surface strength andcompressive residual stress and bending fatigue performance weremeasured and compared with WLSP. The mechanisms of fatigue performanceafter thermal engineered LSP were investigated. The precipitationkinetics as well as dislocation pinning effect affected by precipitateparameters were discussed.

Experiments

Sample Preparation:

The material used in this study was AISI 4140 steel with the followingchemical compositions in wt %: C-0.41, Si-0.21, Mn-0.83, P-0.025,S-0.027, Cr-0.91, Mo-0.18, and Fe-balance. Before LSP, the samples wereaustenitized at 850° C. for 20 min and quenched down to room temperaturein oil. A vacuum furnace was used to carry out tempering treatments byadjusting the tempering time and temperature. The sample dimensions were76.2*10*2.38 mm.

Thermal Engineered LSP Process:

A Q-switched ND-YAG laser (Surelite III from Continuum, Inc.), operatedat the wavelength of 1064 nanometer with a pulse width (full width athalf maximum) of 5 nanoseconds was used to deliver the pulse laserenergy. The laser beam diameter was 1 mm. The laser power intensity wasadjusted by the Q-switched delay time. Aluminum foil with a uniformthickness of 30 microns was applied as the ablative coating. The foilwas firmly attached to the sample surface by the vacuum grease. Insteadof liquid confinements, BK7 glass with a high laser transmissioncoefficient was selected as the confining media. Note that it was morepractical to utilize the silicon oil with a flash point over 600° F. asa confining media for industrial applications of WLSP, since the glasswould likely be shattered by the laser-induced shock pressure. A hotplate was placed below the sample fixture to elevate the workingtemperature, which was monitored by a digital scientific thermometer.After WLSP, the samples were put in a vacuum furnace at controlledtemperature for certain time.

Microstructure Characterization

Transmission electron microscopies (FEI-Tecnai TEM and FEI-Titan TEMoperated at 200 kV and 300 kV respectively) were used to characterizethe microstructures after processing. The focused ion beam (FIB)lift-out method was used to prepare the TEM samples in a FIB instrument(Nova-200). During TEM operation, the Gatan imaging filter (GIF) wasutilized to exclude the inelastic scattered part to improve the qualityof diffraction pattern and observe the weak diffraction spotsoriginating from precipitates.

Mechanical Property Characterization

The surface hardness before and after processing were measured by LecoM-400-H micro-hardness test machine with a 25 gram load and a 10 secondholding time.

The residual stresses were measured by a Bruker D8-Discover X-raydiffraction system using the characteristic X-ray source Co—K_(α). Forthe stress analysis, the peak was used corresponding to a 2θ angle of123.916°, and 11 ψ angles from −50° to 50° were applied and analyzed bythe sin²ψ method. To study the cyclic stability of residual stress andsurface strength, the residual stress and hardness were measured afterdifferent numbers of cyclic loading.

The fatigue properties were determined by the three-point bendingfatigue tests carried out on a 100KN MTS servo-hydraulic fatiguemachine. For all the tests, the load ration, which was defined as theminimum load divided by the maximum load, was set to be 0.10; the samplewas subject to sinusoidally varying load cycles with a frequency of 10Hz at room temperature; the maximal bending stress was calculated byσ=3PL/2bh², where P is the applied load, L is the span for the testset-up, b and h are the width and thickness of specimen.

Results and Discussion

Process Conditions

Warm Laser Shock Peening

Two useful parameters of WLSP process were the processing temperatureand laser power intensity. This was because both DSA and DP effectsduring WLSP increased from the thermal energy provided by the warmprocessing temperature and the high strain rate plastic deformationcaused by the laser-induced shock pressure. The effect of warmprocessing temperature (20° C.≤T_(peen)≤410° C.) during warm shotpeening on mechanical properties of 4140 steel in terms of the halfwidths and surface compressive residual stress have been evaluated. Theoptimal peening temperature was found around 300° C. The optimal WLSPtemperature of 4140 steel in terms of the surface strength wasidentified between 250 and 300° C. A softening effect was observed whilethe peening temperature was higher than 300° C. due to thethermal-induced dynamic recovery. Therefore, to avoid the dynamicrecovery, 250° C. was used as the processing temperature for WLSP inthis study.

To determine the laser intensity for WLSP, surface hardness afterprocessing at 250° C. with laser intensities from 1.2 to 3.6 GW/cm² weremeasured as shown in FIG. 4.1 . It was found the surface hardnessincreased with enhancing the laser power intensity. For instance, thesurface hardness after WLSP increased from 342 to 367VH (kg/mm²) byenhancing the laser intensity from 1.2 to 3.6 GW/cm². This hardeningeffect was attributed to the strain hardening and also the dislocationpinning effect induced by the second phase nano-precipitates. Thegreater laser intensity generated the laser-induced plasma with a higherpeak pressure, which enhanced the strain hardening and results in agreater nucleation rate. The laser-induced peak pressure, P(I₀), as afunction of the applied laser intensity, I₀, was estimated by theFabbro's LSP model as

${{P( I_{0} )} = {0.1( \frac{\alpha}{{2\alpha} + 3} )^{1/2}Z^{1/2}I_{0}^{1/2}}},$where α is the efficiency of the interaction. Z was the combined shockimpedance defined as Z=2/(1/Z₁+1/Z₂), where Z₁ and Z₂ were the shockimpedances of the target material (4140 steel shock impedance 3.96e6 gcm⁻² s⁻¹) and the confining media (BK7 glass shock impedance 1.44e6 gcm⁻² s⁻¹), respectively. The estimated peak pressures for various laserintensities were plotted in FIG. 4.1 . The maximum limit of surfacehardness was observed while the peak pressure was over 4 GPa(corresponding to a laser intensity of 2.4 GW/cm²). This was due to theplastic deformation limit, which was reached when the magnitude of shockpressure was greater than two times of Hugoniot elastic limit (HEL) oftarget metal. This can also be referred to as the plastic deformationsaturation phenomenon. LSP process on three aluminum alloys, gaveplastic deformation reaching a maximum limit while the laser-inducedshock pressure was above two times of HEL. In addition, BK 7 glass wascracked under a high laser power intensity (around 4 GW/cm²). Therefore,to obtain an optimal hardening effect while preventing to crack theconfining media, 3 GW/cm² was utilized as the laser intensity for WLSPin this work.

Pre-Shock Tempering and Post-Shock Tempering

In order to discover the optimal tempering time and study how WLSP wasassisted by tempering, the tempering treatments were carried out beforeand after WLSP, so called pre- and post-shock tempering. The resultingsurface hardness affected by the tempering time was shown in FIG. 4.2 .It was observed that both pre- and post-shock tempered samples couldreach the peak hardness by 2 h tempering treatment. Compared to WLSP, 2h post-tempering could further improve the surface hardness by 28% from360 to 461VHN.

The hardening induced by pre-shock tempering was mainly due to themicrostructure rearrangement, the formation of tempered martensites, andthe diffusion of carbon atoms which were entrapped in the distortedbody-centered-tetragonal (BCT) structure through quenching.

For the post-shock tempered samples, the hardening was attributed to thedislocation-assisted coarsening of nano-precipitates. Additionally, witha longer tempering time (longer than 2 hours), a helpful drop of surfacehardness in post-shock tempered samples was seen. This was caused by theover-tempering effect, which dramatically decreased the precipitatenumber density and resulted in the thermal-induced reduction ofdislocation density. The tempering effect on microstructure evolutionand mechanical properties studied with details in the followingsections. Furthermore, the magnitude of peak hardness of post-shocktempered sample (461VHN) was greater than that of pre-shock temperedones (408VHN). Since the resistance to crack initiation and propagationwas affected by the surface strength, 2 h post-shock tempering wasselected as the optimal tempering condition for the fatigue performance.

Post-Shock Tempering Temperature

The post-shock tempering temperature was another useful processingparameter in thermal engineered LSP process. FIG. 4.3 shows the surfacehardness affected by post-shock tempering temperature at varioustempering time. A noticeable hardness drop was observed in the samplestempered at 550° C. This softening effect was attributed to thethermal-induced dynamic recovery, which leads to a helpful reduction ofdislocation density. Favorable results were realized utilizing temperingtemperatures from approximately 350° C. to approximately 450° C.However, compared to the samples tempered at 350° C., those tempered at450° C. could reach the peak hardness with a greater magnitude in ashorter tempering time. For instance, the peak hardness of 450° C.tempered samples reached by 2 h post-shock tempering is 461VHN, whilethat of 350° C. tempered sample was 436VHN but requiring 4 h temperingtreatment. In addition, the possibility of tempered martensiteembrittlement in 4140 steel within the temperature range between 250 and400° C. has was reported, which could accelerate the crack initiation.Therefore, to avoid the thermal recovery and obtain the peak hardness,450° C. was used as the tempering temperature for thermal engineered LSPof 4140 steel.

Microstructures Beneath the Peened Surface

The mechanical properties of metallic materials were governed by themicrostructure behaviors, such as the dislocations' collective motionand their interactions among themselves and with other crystal defects.It was useful to investigate the structural evolutions taking placeduring thermal engineered LSP at the zone directly beneath the peenedsurface. This compressive zone extends to a length scale of severalhundred micrometers, which was significantly larger than the deformedlayer. TEM images in FIG. 4.4 showed the initial microstructures ofoil-quenched sample, in which dislocation structures with a low densitywere observed. The generation of these dislocations was majorlyattributed to the crystal structure transition of steel from theface-centered cubic (austenite phase) to the body-centered tetragonal(martensite phase). The microstructure evolution in thermal engineeredLSP process was analyzed by considering three processing stages: WLSP,the short time post-shock tempering and over-tempering.

Microstructures Generated by WLSP

FIG. 4.5 depicts weak beam dark field TEM images (FIGS. 4.5 a and 4.5 b) showing the entanglement of dislocations and spherical-shapednano-precipitates in WLSP sample and the related diffraction pattern oftwo beam condition (FIG. 4.5 c ). It was discovered that the tangleddislocation structures with a high density and a uniformly distributedarrangement were formed rather than the dislocation pile-ups or shearbands, which were normally formed in metallic materials after to roomtemperature LSP due to the high strain rate deformation. This was mainlycaused by the enhanced dislocation multiplication due to DSA. In WLSP,the warm processing temperature provides sufficient thermal energy forthe diffusion of solute atoms (carbon) to dislocation cores. Themigration of solute atoms results in the formation of Cottrellatmosphere, which exert a pinning force to resist the movement of nearbymobile dislocations. To continue the plastic deformation, the greaterstress was useful to overcome the local resistance for dislocationmotion. This led to the enhanced dislocation multiplication byactivating more dislocation sources and generating new mobiledislocations.

The uniformly distributed dislocations resulted from DSA and also thepinning effect of highly dense nano-precipitates. Thesenano-precipitates (pointed out by arrows in FIG. 4.5 a ) were generatedsimultaneously during the laser-induced plastic deformation through DP.In WLSP, the high strain rate plastic deformation results in theformation of highly dense dislocation nodes, which were generallyconsidered as the potential nucleation sites for precipitation. Both thewarm processing temperature and high strain rate plastic deformationassist and accelerate nucleation growth during WLSP by decreasing theactivation energy for precipitation. The details of nucleation kineticsaffected by the processing temperature and laser shock are discussed inthe precipitation kinetics section. Highly dense nano-precipitates weregenerated by WLSP due to DP. The size of these particles is around 5-10nm in diameter, which were interpreted by the short aging time (in thescale of nano-second).

Microstructures Generated by Post-Shock Tempering

After WLSP, as found by the surface hardness test (FIG. 4.3 ), a shorttime post-shock tempering treatment (shorter than 2 hours) introduced afurther hardening effect on the surface strength; on the other hand, asoftening effect resulted from the long time post-shock temperingtreatment (longer than 4 hours). It was helpful to understand thishardening and softening mechanism from the point of view ofmicrostructure evolution during post-shock tempering.

The TEM images in FIG. 4.6 and FIG. 4.7 compare the microstructuresinduced by the short time (2 hours, FIG. 4.6 ) and the long time (6hours, FIG. 4.7 ) post-shock tempering treatments in terms ofdislocation and precipitate structures. It was observed that thedislocation structures in 2 h-tempered sample remain the relativelyuniformly distributed arrangement, which was introduced through DSAduring WLSP. The pile-ups of localized dislocations with a high densitywere formed in the 6 h-tempered sample. These dislocation pile-ups werealso called dislocation walls (indicated by arrows in FIGS. 4.7 a and4.7 b ), which were generated due to the thermal-induced recovery ofdislocations. This microstructure recovery could be identified in 2h-tempered sample due to the presence of highly dense fine precipitates,which act as local obstacles interact elastically with nearby movingdislocations. The pinning effect made a sizeable contribution to resistthe dislocation motion and retard the recrystallization normally inducedduring the heat treatment after deformation. In this way, the presenceof precipitates helps to retain the deformed structure and accumulatedstrain. However, a long time post-shock tempering treatment induced thecoarsening of precipitates with a helpful decrease in precipitate numberdensity. This made the pinning effect less effective, and hence lead tothe formation of dislocation walls and reduction of dislocation density.A similar phenomenon was observed in previous work after a short timepost-deformation aging at the temperature up to 600° C., theequal-channel angular pressing (ECAP) structure of Cu—Cr—Zr alloyremains fine-grained and dislocation density remains high due to theimproved thermal stability resulted from the pinning effect ofnano-grains and nano-precipitates.

To have a better sight on the precipitate structures, a selected areadiffraction (SAD) technique was applied to obtain the weak diffractionspots associated with the precipitate structures in the diffractionpattern. The dark field TEM images of typical precipitate structures in2 h-tempered and 6 h-tempered samples were shown in FIGS. 4.6 c, 4.6 d,4.7 c and 4.7 d . The selected weak spots were circled and indicated indiffraction patterns. The coarsening of precipitates during post-shocktempering was observed. For instance, the average size of precipitatesin 6 h-tempered sample was around 100 nm in diameter, which was largerthan that of 2 h-tempered sample (40-50 nm in diameter) and WLSP sample(5-10 nm in diameter). However, this coarsening phenomenon wasaccompanied with a helpful decrease of precipitate number density ascompared by FIG. 4.6 and FIG. 4.7 . In addition, it was discovered thatthe coarsening of precipitates takes place at the same region wherelocalized dislocation pile-ups were formed. As observed from FIGS. 4.7 aand 4.7 b , the large particles (circled by rings) were mostly foundalong dislocation wall structures (indicated by arrows). This suggests amutual influence of precipitates and dislocations: dislocation pile-upsprovided favorable nucleation sites for the nucleation and coarsening ofprecipitates, while dislocation slips were restrained by the presence ofprecipitates due to the pinning effect. The HRTEM images of precipitatesfor 2 h post-shock tempered and 6 h post-shock tempered samples areshown in FIG. 4.6 a and FIG. 4.8 b

Coarsening of Precipitates

FIG. 4.8 shows the evolution of precipitate radius R, number density Nand volume fraction f_(v) with the post-shock tempering time. The dataof number density were estimated by TEM images taken from the regionclose to the sample surface. For the WLSP sample, in which the fineparticles could not be identified due to the small size, the numberdensity was referred from the reported WLSP nucleation model. The volumefraction was calculated by f_(v)=4πR³N/3. The results showed that thecoarsening of precipitates with post-tempering time was accompanied bythe linear decrease of particle number density, which was accordancewith the Lifshitz-Slyozov-Wagner (LSW) theory (n∝t⁻¹). Furthermore, arapid increase in volume fraction at the beginning of tempering processwas observed, followed by a slowdown at longer tempering time. Thesimilar phenomenon was reported in a study of the behaviors of thehardening precipitates in Mg—Y—Nd alloy during aging treatments. It wasfound that at the early stage of aging treatment, a rapid increase ofprecipitate volume fraction was accompanied by a linear decrease ofparticle number density.

Precipitation Kinetics

In the nucleation stage, as a dynamic precipitation process, thenucleation of precipitates was dislocation-assisted, and the dislocationnodes formed by deformation were favored nucleation sites forprecipitation. The available nucleation site density N₀, employed todescribe the sites where it was possible for the nucleation to takeplace, was one of the parameters for the nucleation kinetics. With theconsideration of laser intensity and processing temperature effects, theavailable nucleation site density could be expressed as:

$\begin{matrix}{{N_{0}( {I_{0},T} )} = {{\frac{1}{2}{\rho( {I_{0},T} )}^{3/2}} = ( \frac{{\sigma( {T,I_{0}} )} - \sigma_{0}}{M\gamma\mu b} )^{2}}} & (1)\end{matrix}$where M was the Taylor factor (≈3.10), μ and b are the shear modulus andBurgers vector of dislocations, γ was a scalar coefficient between 0.2and 0.4, and σ was the flow stress of metallic materials afterprocessing affected by laser intensity I₀ and processing temperature T,and σ₀ was the intrinsic friction stress contributing to the flowstress.

Besides the available nucleation site density, the nucleation rate wasalso determined by the nucleation coefficient, which was affected by theprocessing temperature and the activation energy for nucleation.Accordingly, the nucleation rate could be obtained by:

$\begin{matrix}{\frac{\mathbb{d}N}{\mathbb{d}t}{{\lambda(T)} \cdot N_{0}}} & (2)\end{matrix}$ $\begin{matrix}{{\lambda(T)} = {Z^{\prime}\beta^{\prime}{\exp( {- \frac{\Delta G}{kT}} )}{\exp( {- \frac{\tau}{t}} )}}} & (3)\end{matrix}$where N was the number of precipitates per unit volume, λ(T) is thenucleation coefficient (between 0 to 1) affected by the processingtemperature, Z′ was Zeldovich's factor, t and β′ was the time and atomicimpingement rate, and r was the incubation period. The activation energyfor nucleation ΔG was consisted of four components:ΔG=V(ΔG _(chem) +ΔG _(c))+γ_(i) −ΔG _(d)  (4)where V is the volume of nucleus, ΔG_(chem), ΔG_(E), γ_(h) and ΔG_(d)are the chemical driving force, volume strain energy, interfacial freeenergy and dislocation core energy relatively.

With utilizing of the extended mechanical threshold stress (MTS) modelfor flow stress (feasible for the ultra high strain rate deformationgreater than 10⁵/s) and Fabbro's LSP model for laser shockwave peakpressure, the nucleation rate during WLSP was numerically calculated andverified by experiments. Both calculation results and experimentalobservations confirm that the nucleation kinetics was assisted andaccelerated by the deformation-induced dislocations and the elevatedprocessing temperature. This could potentially be understood by thefollowing reasons from the activation energy point of view. First ofall, the activation energy for nucleation was reduced by the increaseddislocation core energy caused by the stronger dislocation interactions.The dislocation core energy stored in dislocation nodes during themechanical deformation was released for the nucleation growth in theprecipitation process. With a higher dislocation density induced by thelaser shock, a greater dislocation core energy is stored in dislocationnodes, resulting in a lower requirement of external applied energy fornucleation. Furthermore, the activation energy for nucleation wasreduced by the presence of highly dense dislocations due to the reducedvolume strain energy. The volume strain energy related to thematrix/precipitate lattice mismatch was proportional to the elasticlattice dis-registry and inversely proportional to the coherency lossparameter. While the precipitates were surrounded by highly densedislocations, the matrix/precipitate interface was semi-coherency, whichmeans a lower elastic lattice dis-registry and a greater value ofcoherency loss parameter. This leads to the reduced volume strain energyfor nucleation. In addition, the activation energy for nucleation wasreduced by the elevated processing temperature due to the lower chemicaldriving force. The chemical driving force represents the free energydifference between matrix and precipitates. Before reaching theequilibrium temperature (at which the free energy of matrix andprecipitate phase were equal), the excess free energy provided by matrix(unstable phase) relative to the precipitate (stable phase) was reducedby increasing temperature.

In the coarsening stage, the coarsening apparently took place at thedislocation nodes, since the precipitates were interconnected throughthe dislocation network. The effective diffusion coefficient wascontributed from the bulk diffusion and also the pipe diffusion throughdislocations. The effective diffusivity D_(e) was expressed as:D _(e) =D _(b)+(D _(p) −D _(b))πR _(core) ²ρ  (5)where D_(b) was the bulk diffusivity, D_(p) was the pipe diffusioncoefficient, R_(core) was the radius of the dislocation core, and ρ wasthe total dislocation density. Based on the classic precipitationtheory, the coarsening rate expressed by precipitate radius and numberdensity evolutions were obtained as:

$\begin{matrix}{\frac{\mathbb{d}R}{\mathbb{d}t_{coarsening}} = \frac{4C^{e}R_{0}D_{e}}{{27C^{p}} - {C^{e}R^{2}}}} & (6)\end{matrix}$ $\begin{matrix}{\frac{\mathbb{d}R}{\mathbb{d}t_{coarsening}} = {\frac{4C^{e}R_{0}D_{e}}{{27C^{p}} - {C^{e}R^{3}}}\lbrack {{\frac{R_{0}C}{R( {C^{p} - C} )}( {\frac{3}{4\pi R^{3}} - N} )} - {3N}} \rbrack}} & (7)\end{matrix}$where C^(e) was the equilibrium concentrations of carbon at theannealing temperature, C^(p) was the concentration of carbon inprecipitate, and C was the solute concentration of carbon. The effect ofprior deformation on precipitation kinetics at the coarsening stage wasnumerically investigated taking care of both the bulk diffusion and thepipe diffusion. It was found that, for the prior deformed0.084C-0.015N-0.06Nb steel sample subjected to the annealing treatment,the growth of precipitate size was more helpful in the long annealingtime regime than that in the short time regime, and the precipitatenumber density increased rapidly at the early stage of annealing,followed by a notable decrease at the longer annealing time. The similarcoarsening phenomenon is observed by TEM images before. In-depth studyof the coarsening mechanism during the post-shock tempering treatmentplayed a useful role for theoretically understanding the microstructureevolution during thermal engineered LSP process.

Effect of Dislocation Pinning on Mechanical Performances

Dislocation Pinning Strength Affected by Particle Parameters

Strengthening due to second phase precipitates has been examined. At lowtemperatures, dislocation overcomes local obstacles by cutting throughor looping around mechanisms. The former process took placepredominantly for small precipitates with low density, while the latterfor large precipitates with high density. In the case of cuttingmechanism, the Friedel model can be applied to estimate the criticalshear stress τ_(c), which could be expressed as

$\begin{matrix}{\tau_{c} = {C\frac{{F_{\max}^{3/2}({DN})}^{1/2}}{{b( {12S} )}^{1/2}}}} & (8)\end{matrix}$where S was the linear tension of the dislocation, b was the magnitudeof the Burgers vector, D and N were the size and volume number densityof precipitates, C˜1 was a constant, and F_(max) was the maximum valueof dislocation/particle interaction force. Eq. 8 suggests that thecritical shear stress with a greater magnitude can be useful for thedislocation movement in the metal matrix containing precipitates with ahigher number density and larger sizes. For small particles, thestrengthening induced by cutting may be contributed by at least twomechanisms: (1) the stress fields induced by particles, and (2) themodulus interaction between dislocation and particle, which resulted inthat the maximum force proportional to the difference in the elasticmoduli of matrix and precipitate (F_(max)˜Δμ/μ). With the increase ofparticle size, the dislocation appeared to be pinned inside theprecipitate while cutting through, resulting in a notable increase ofthe critical shear stress as compared to that induced by modulusinteraction. This lead to the changeover of dislocation/precipitateinteraction mechanism from cutting through to Orowan looping, which leftobstacles surrounded by pinned dislocation loops after bypassing. Thestrengthening contribution of particles in looping around process wasmainly considered by the critical resolved shear stress τ_(p) (theminimal external stress for overcoming dislocation/precipitateinteraction force) as a function of the mean particle radius r andvolume fraction f. The advanced τ_(p) function based on the worked wasas follows, which takes into account the actual distribution of particlesizes.

$\begin{matrix}{{{\tau_{p}( {r,f} )} = {\frac{0.9b\mu}{2\pi w_{L}r\sqrt{1 - v}}{{\ln( \frac{2w_{D}r}{b} )}\lbrack \frac{\ln( {2w_{D}{r/b}} )}{\ln( {w_{L}{r/b}} )} \rbrack}^{1/2}}}{with}} & \end{matrix}$ $\begin{matrix}{{w_{L} = {\sqrt{\pi{w_{q}/f}} - {2w}}},{w_{D}^{- 1} = {w_{L}^{- 1} + ( {2w_{r}} )^{- 1}}}} & (9)\end{matrix}$where μ and v were the shear modulus and Poisson ratio of matrix, w_(r)and w_(q) were the mean radius and the mean area of the particleintersection with the glide plane. In the case of Orowan looping, agreater critical resolved shear stress was required with the increase ofparticle size and volume fraction. With a larger particle size, theparticle-induced stress field was enhanced, and the bigger dislocationloops had to be formed after bypassing; with a higher particle volumefraction, dislocations had to bend more strongly due to the less freespace between particles. Therefore, a higher external stress wasrequired for overcoming the dislocation/precipitate interaction. Amulti-scale discrete dislocation dynamic simulation was utilized toinvestigate the particle/dislocation interaction during laser peeningprocess. The dislocation pinning effect on the microstructure evolutionand material stress-strain curves were studied by initially insertingprecipitates with various sizes, densities and volume fractions into thecomputational cells. With the consideration of precipitate volumefraction, the optimal pinning strength was determined by the balancebetween the particle size and number density effects. It was worthwhileutilizing thermal engineered LSP experiments to verify the simulatedstrengthening effect by considering the surface hardness (FIG. 4.2 ) andalso the cyclic stability of surface strength and residual stress.

Cyclic Stability of Surface Strength and Residual Stress

Besides the magnitude and depth, the stability of compressive residualstress played another role in determining the fatigue performance ofmetallic materials. FIG. 4.9 compares the residual stress relaxationduring the cyclic loading at 1400 MPa maximal three-point bending stressof WLSP and 2 h post-shock tempered samples. Compared to WLSP sampleswith no additional tempering, the initial residual stress before cyclicloading of 2 h post-shock tempered samples had a lower magnitude. Thisstress relaxation was mainly caused by the thermal-induced diffusionalcreep and dislocation glide during the tempering process. However, alarge stress relaxation was seen in WLSP samples with increasing cyclicloading numbers. For example, after 100 k cyclic loading the residualstress of WLSP samples decreased by 43.9% from 421 to 236 MPa, whilethat of 2 h tempered samples decreased by 21.6% from 357 to 280 MPa.These results demonstrated that a higher cyclic stability of residualstress was achieved by 2 h post-shock tempering treatment. In WLSPsamples, the cyclic residual stress relaxation was mainly due to themicro-plastic strains accumulating from cycle to cycle, while in 2 htempered samples, the presence of larger size precipitate couldeffectively resist material's plastic behaviors by locking mobiledislocations. In addition, the stress relaxation occurred when thesuperposition of the loading stress, residual stress and dislocationpinning stress exceeds material's yield stress. Therefore, for 2 htempered sample, with the same loading stress and a lower initialresidual stress, the suppressed stress relaxation was attributed to theenhanced pinning stress. This improvement makes a significantcontribution for a better fatigue performance.

The stability of surface strength was also useful for the effectivefatigue improvements by determining the resistance to crack initiation.The surface hardness changed with the cyclic loading of samplesprocessed by various conditions was shown in FIG. 4.10 . It was ofinterest to notice that in the first 100 cycles, a helpful hardeningeffect was induced by cyclic loading in 2 h tempered samples, while nosuch phenomenon was observed in WLSP samples. It was generally acceptedthat the cyclic hardening was attributed to the dislocationmultiplication, dislocation mutual interaction and/or theprecipitate/dislocation interaction. Both WLSP and 2 h tempered samplescontain the highly dense dislocations, but the precipitate size of 2 htempered sample was 4 times larger than that of WLSP sample. Therefore,the noticeable cyclic hardening in 2 h tempered sample can be explainedby the enhanced pinning effect from the precipitate coarsening. Theresulted larger particles exert a greater pinning force to resist themovement of dislocations during the cyclic loading. To continue theplastic behavior, the generation of more mobile dislocations wasnecessary, resulting in the enhanced dislocation multiplication. Thiscontributed to a greater surface hardness. After 100 cycle loadings, asoftening effect of surface hardness was observed with increasing cyclenumbers in both WLSP and 2 h tempered samples. However, compared to WLSPsample, the 2 h tempered sample had a lower softening rate, whichindicated a higher cyclic stability of surface strength. For example,after 100K cycle loadings, the hardness of WLSP sample reduced by 33.1%from 344 to 230VHN, while that of 2 h-post tempered sample only had a4.4% reduction from 452 to 432VHN. This cyclic softening effect could beexplained by the annihilation of dislocations or the rearrangement ofdislocations to subgrains. Due to the presence of larger sizeprecipitates, the movement and rearrangement of dislocations is resistedby the pinning strength, resulting in the suppressed cyclic softeningeffect. Therefore, a higher cyclic stability of surface hardness wasachieved by 2 h post-shock tempering treatment, which aspects to thefatigue performance through the resistance to crack initiation andpropagation.

Fatigue Test

The three-point bending fatigue performances that are obtained aftervarious processing conditions are depicted in the stress-lifespan (S—N)curves in FIG. 4.11 . It was shown that compared to WLSP, a betterfatigue performance is achieved by 2 h post-shock tempering treatment.For example, under the certain stress magnitude from 1500 MPa to 1750MPa, the fatigue life of 2 h tempered specimens are 2-10 times higherthan that of WLSP specimens. In addition, the bending fatigue strengthof 4140 steel after 2 h post-shock tempering was enhanced by 200 MPafrom 1300 MPa to 1500 MPa. This helpful improvement was 2 times greaterthan that induced by WLSP relative to LSP. Note that, compared to LSP,WLSP improves the bending fatigue strength of 4140 steel by 75 MPa. Theimproved fatigue strength was of interest because it is not in keepingwith the commonly held concept that the specimen after WLSP with ahigher magnitude of compressive residual stress (as shown in FIG. 4.9 )should exhibit a better fatigue performance, and 2 h post-shocktempering treatment should have a negative effect on fatigue strengthdue to the thermal-induced stress relaxation. According to the previousTEM images and cyclic stability study, the improvement of fatigueperformance appears to be substantially due to the enhanced dislocationpinning effect resulted from the coarsening of precipitates duringtempering. The resulting larger size precipitates exert a greaterpinning force to stabilize dislocation structures during cyclic loading.As a result, the applied stress with higher amplitude was needed toovercome the pinning effect to induce the movement of dislocations.Meanwhile, the crack initiation was resisted by the greater surfacehardening, and the crack propagation was suppressed by the compressiveresidual stress with higher cyclic stability. Therefore, a betterfatigue performance was achieved by 2 h post-shock tempering. On theother hand, a longer time post-shock tempering had a negative effect onthe fatigue performance. As observed from FIG. 4.11 , the S—N curve of 4h tempered sample moves to the left of that of 2 h tempered sample.Additionally, compared to 2 h tempered specimens, the fatigue strengthafter 4 h post-shock tempering was decreased by 100 MPa. This negativeeffect is attributed to the helpful reduction of particle number densityduring longer time tempering treatment, which lead to a softening effecton the pinning strength.

Therefore, with carefully manipulated processing conditions, the optimalmicrostructures with the optimized pinning effect of metallic materialscould be reached by thermal engineered LSP for mechanical applicationsincluding enhancing the surface hardness, improving the cyclic stabilityof surface strength and residual stress, and extending the fatigue life.Compared to other surface processing techniques like shot peening, deeprolling, LSP etc, thermal engineered LSP has a higher potential to reachthe optimized mechanical properties, since this technique integrates thefollowing advantages: (1) the effective strain hardening effect and thehigh magnitude and in-depth compressive residual stress produced duringLSP; (2) dynamic strain aging for the highly dense dislocations with auniform arrangement; (3) dynamic precipitation for the high nucleationrate of nano-precipitates; and (4) the post-shock tempering for theoptimized pinning strength. Thermal engineered LSP should have manybeneficial industrial applications.

Thermal Engineered Laser Shock Peening Conclusion

Thermal engineered LSP was carried out on AISI 4140, and the resultingmechanical improvements were evaluated. The mechanisms of fatigueperformance after thermal engineered LSP were investigated. Thermalengineered LSP combines WLSP with the followed post-shock temperingtreatment to extend the precipitation kinetics from the nucleation stageto the coarsening stage. In the nucleation stage, both DP and DSAeffects lead to the formation of highly dense nano-precipitatessurrounded by highly dense dislocations. In the coarsening stage, withthe assistance of dense dislocations, the precipitate size grows larger,while the number density is decreased. By carefully manipulatingtempering conditions, the precipitate size and number density effects onthe dislocation pinning strength could be balanced to obtain theoptimized pinning strength, which effectively contributes to stabilizemicrostructures by locking mobile dislocations during cyclic loading. Asa result, the cyclic stability of surface strength and compressiveresidual stress induced by laser shock is improved, and thus the fatiguelife is extended. According to mechanical test results, 2 h post-shocktempering at 450° C. after WLSP leads to the optimized microstructure ofAISI 4140. Compared to WLSP, 2 h post-shock tempering could furtherenhance the surface hardness by 28%, and improve the fatigue limit by200 MPa.

Various aspects of different embodiments of the present disclosure areexpressed in paragraphs X1, X2 and X3 as follows:

X1. One embodiment of the present disclosure includes an apparatus forhardening a material, comprising: a heater that heats the material to adesired temperature; and a laser oriented to emit a beam of laser energytoward the material; wherein the laser generates a plasma proximate thematerial and deforms the surface of the heated material.

X2. Another embodiment of the present disclosure includes an apparatusfor hardening a material, comprising: a cooling member that cools thematerial to a desired temperature; and a laser oriented to emit a beamof laser energy toward the material; wherein the laser generates aplasma proximate the material and deforms the surface of the cooledmaterial.

X3. A further embodiment of the present disclosure includes an apparatusfor hardening a material, comprising: a container adapted to contain aliquid active agent as the agent covers the material; and a laseroriented to emit a beam of laser energy through the active agent;wherein the laser decomposes at least a portion of the active agent,creates a shock, and deforms the surface of the material

Yet other embodiments include the features described in any of theprevious statements X1, X2 or X3, as combined with one or more of thefollowing aspects:

Wherein the heater maintains the desired temperature at a temperaturethat maximizes the density of nanoprecipitation.

Wherein the desired temperature is at least 20 deg. C. and at most 400deg. C.

Wherein the material is an aluminum alloy and the desired temperature isat least approximately 150 deg. C. and at most approximately 200 deg. C.

Wherein the material is steel and the desired temperature is at leastapproximately 250 deg. C. and at most approximately 400 deg. C.

Wherein the temperature is maintained within 20 deg. C. of the desiredtemperature.

Wherein the temperature is maintained within 20 deg. K of the desiredtemperature.

Wherein the heater heat tempers the material after said deforming.

Wherein the heat tempering includes heating the material to at least 250deg. C. and at most 450 deg. C. for at least 2 hours.

Wherein the heat tempering includes heating the material toapproximately 450 deg. C. for approximately 2 hours.

Wherein the heater heat tempers the material prior to said deforming.

Wherein nanomaterials are generated.

Wherein the desired temperature is selected to maximize mechanicaltwining in the material.

Wherein the desired temperature is at least 50 deg. K and at most 250deg. K.

Wherein the active agent is in direct physical contact with thematerial.

Wherein the material is covered with an ablative member positionedbetween the material and the active agent.

Wherein a plasma is generated with the laser energy.

Wherein the active agent is a liquid during said emitting.

Wherein the active agent is hydrogen peroxide.

Reference systems used herein may refer generally to various directions(e.g., upper, lower, forward and rearward), which are merely offered toassist the reader in understanding the various embodiments of thedisclosure and are not to be interpreted as limiting. Other referencesystems may be used to describe various embodiments, such as referringto the direction of projectile movement as it exits the firearm as beingup, down, rearward or any other direction.

While examples, one or more representative embodiments and specificforms of the disclosure have been illustrated and described in detail inthe drawings and foregoing description, the same is to be considered asillustrative and not restrictive or limiting. The description ofparticular features in one embodiment does not imply that thoseparticular features are necessarily limited to that one embodiment.Features of one embodiment may be used in combination with features ofother embodiments as would be understood by one of ordinary skill in theart, whether or not explicitly described as such. One or more exemplaryembodiments have been shown and described, and all changes andmodifications that come within the spirit of the disclosure are desiredto be protected.

What is claimed is:
 1. A method for hardening a material, comprising:selecting a desired temperature to maximize a density ofnanoprecipitation in a material, wherein the desired temperature is lessthan a recrystallization temperature of the material and greater thanroom temperature; heating, using a heater, the material to the desiredtemperature, wherein the desired temperature is at least 20 deg. C. andat most 400 deg. C., and maintaining the desired temperature, using theheater; emitting a pulse of laser energy toward the material, thematerial being maintained at the desired temperature, the pulse having aduration of less than 100 nanoseconds; generating a plasma with thelaser; confining the plasma with a confining material, wherein theconfining material includes fused quartz or silicone oil, and whereinexpansion of the confined plasma generates a shock wave proximate thesurface of the material; deforming a surface of the material as a resultof the shockwave; and generating nanoprecipitates internal to saidmaterial by said deforming at the desired temperature; wherein theheater is distinct from the source of the pulse of laser energy.
 2. Themethod of claim 1, wherein the material is an aluminum alloy and thedesired temperature is at least approximately 150 deg. C. and at mostapproximately 200 deg. C.
 3. The method of claim 1, wherein the materialis steel and the desired temperature is at least approximately 250 deg.C. and at most approximately 400 deg. C.
 4. The method of claim 1,comprising: maintaining the material at a temperature within 20 deg. C.of the desired temperature.
 5. The method of claim 1, comprising: heattempering the material after said deforming.
 6. The method of claim 5,wherein said heat tempering includes heating the material to at least250 deg. C. and at most 450 deg. C. for at least 2 hours.
 7. The methodof claim 5, wherein said heat tempering includes heating the material toapproximately 450 deg. C. for approximately 2 hours.
 8. The method ofclaim 1, comprising: heat tempering the material prior to saiddeforming.
 9. The method of claim 1, comprising: emitting a plurality ofpulses of laser energy toward the material, the duration of each pulsebeing less than 100 nanoseconds.
 10. The method of claim 1, comprising:coating the surface of the material with an ablative coating before saidemitting.
 11. The method of claim 10, wherein the ablative coating isselected from the group consisting of: metallic foil, organic paint andadhesive.
 12. The method of claim 10, wherein the ablative coatingincludes metallic foil, organic paint or adhesive.
 13. The method ofclaim 10, which further comprises confining the surface of the materialbefore said emitting.
 14. The method of claim 1, wherein said deformingthe surface is plastically deforming the surface.
 15. The method ofclaim 14, wherein said deforming the surface is near-surfacedeformation.
 16. The method of claim 1, which further comprisesdynamically precipitating a microstructure of the material by saiddeforming.
 17. The method of claim 1, which further comprisesdynamically strain aging a microstructure of the material by saiddeforming.
 18. The method of claim 1, which further comprises hardeningthe surface of the material as a result of said deforming.
 19. Themethod of claim 1, which further comprises confining the surface of thematerial before said emitting.
 20. The method of claim 1, which furthercomprises hardening the surface of the material as a result of saiddeforming.
 21. The method of claim 1, wherein the confining material issilicone oil.